EP0020505B2 - Verfahren zur herstellung von aluminium-legierungen - Google Patents

Verfahren zur herstellung von aluminium-legierungen Download PDF

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Publication number
EP0020505B2
EP0020505B2 EP79901364A EP79901364A EP0020505B2 EP 0020505 B2 EP0020505 B2 EP 0020505B2 EP 79901364 A EP79901364 A EP 79901364A EP 79901364 A EP79901364 A EP 79901364A EP 0020505 B2 EP0020505 B2 EP 0020505B2
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Prior art keywords
alloy
strength
plate product
fracture toughness
present
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French (fr)
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EP0020505A4 (de
EP0020505A1 (de
EP0020505B1 (de
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William E. Quist
Michael V. Hyatt
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Boeing Co
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Boeing Co
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C22C21/10—Alloys based on aluminium with zinc as the next major constituent
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/04—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of aluminium or alloys based thereon
    • C22F1/053—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of aluminium or alloys based thereon of alloys with zinc as the next major constituent

Definitions

  • the present invention relates to a method of producing aluminum alloys, and more particularly to a method of producing alloys of 7000 series of the aluminum-zinc-magnesium-copper type characterized by high strength, high fatigue properties and high fracture toughness.
  • alloy 7075 in the T651 temper.
  • Alloy 7075-T651 has a high strength to weight ratio, while exhibiting good fracture toughness, good fatigue properties, and adequate corrosion resistance.
  • Another currently available alloy sometimes used on commercial jet aircraft alloy 7178-T651 is stronger than 7075-T651; however, alloy 7178-T651 is inferior to alloy 7075-T651 in fracture toughness and fatigue resistance.
  • alloy 7075-T651 Other currently available alloys and tempers, although sometimes exhibiting good toughness properties and high resistance to stress-corrosion craking and exfoliation corrosion, offer no strength advantage over alloy 7075-T651. Examples of such alloys are 7475-T651, T7651 and T7351 and 7050-T7651 and T73651. Thus with currently available alloys and tempers, it is impossible to achieve a weight saving in aircraft structural components while maintaining fracture toughness, fatigue resistance and corrosion resistance at or above the level currently available with alloy 7075-T651.
  • the 7000 series alloy produced according to the present invention fulfills the foregoing objects by providing a strength increase of from 10 to 15% over alloy 7075 in T6 tempers. Indeed, the alloy produced according to the present invention is stronger than any other commercially available aluminum alloy. At the same time, the fracture toughness and fatigue resistance of the aluminum alloy produced according to the present invention are higher than that achievable in alloys having strengths approaching that of the alloy of the present invention, such as 7075 and 7178 in the T6 tempers. Additionally, the corrosion resistance of the alloy produced according to the present invention is approximately equivalent to that exhibited by alloy 7075 in the T6 tempers.
  • the desired combination of properties of the aluminum alloy produced according to the present invention has been achieved in a 7000 series alloy by precisely controlling the chemical composition ranges of the alloying and trace elements, by heat treating the alloy to increase its strength to high levels, and by maintaining a substantially unrecrystallized microstructure.
  • the invention relates to a method of producing an aluminum plate product for an upper wing skin for an aircraft, said method comprising the steps of:
  • the high strength, high fatigue resistances, high fracture toughness and corrosion resistance properties of the alloy produced according to the present invention are dependent upon a chemical composition that is closely controlled within specific limits as set forth below, a carefully controlled heat treatment of products made from the alloy, and a microstructure that is substantially unrecrystallized. If the composition, fabrication, and heat treatment parameters of the invention alloy stray from the limits set forth below, the desired combination of strength increase, fracture toughness increase and fatigue improvement objectives will not be achieved.
  • the aluminum alloy produced according to the present invention consists of 5.9 to 6.9% zinc, 2.0 to 2.7% magnesium, 1.9 to 2.5% copper, 0.08 to 0.15% zirconium, the balance being aluminum and trace elements.
  • the maximum percentage of iron allowable is 0.15%
  • of silicon allowable is 0.12%
  • of manganese allowable is 0.10%
  • of chromium allowable is 0.04%
  • titanium allowable is 0.06%.
  • Any other remaining trace elements have maximum limits of 0.05%, with a maximum total for the remaining trace elements being 0.15%. (The foregoing percentages are weight percentages based on the total alloy).
  • the most critical of the trace elements present are normally iron and silicon. If the iron and silicon are present in the alloy in excess of the amounts stated above, the undesirable intermetallic compounds formed by iron and silicon during solidification, fabrication, and heat treatment will lower the fracture toughness properties of the alloy of the present invention to unacceptable levels.
  • the high zinc, magnesium and copper contents of the alloy produced according to the present invention are major contributors to the high strength characteristics of the present alloy. If the zinc, magnesium and copper contents are below the limits set forth above, the strength of the alloy will fall below the strength objectives of 10-15% increase over that of the base line standards, alloy 7075-T651.
  • Conventional melting and casting procedures are employed to formulate the alloy. Care must be taken, as pointed out above, to maintain high purity in the aluminum and the alloying constituents so that the trace elements, and especially iron and silicon, are maintained below the requisite maximums.
  • Ingots are produced from the alloy using conventional procedures such as continuous direct chill casting. Once the ingot is formed, it can be homogenized by conventional techniques, for example, subjecting the ingot to elevated temperatures of about 482°C.
  • the product formed from an alloy of the present invention must be substantially unrecrystallized.
  • substantially unrecrystallized it is meant that less than about 50 volume percent of the alloy microstructure in a given product is in a recrystallized form, excepting surface layers which often show a much higher degree of recrystallization. (The surface layers of plate and products are usually removed during fabrication into final part configurations). Most preferably, it is desired to maintain the volume percent of recrystallized microstructure less than about 30%. Recrystallization can be minimized by maintaining the temperature during hot rolling at levels that cause annealing out of internal strains produced by the rolling operations such that recrystallization will be minimized during the rolling operation itself, or during subsequent solution treatment.
  • the product is typically solution heat treated at a temperature on the order of 477°C., and preferably between 477°C. and 482°C. for a time sufficient for solution effects to approach equilibrium.
  • the product is quenched, normally by spraying the product with, or immersing the product in, room temperature water. Thereafter the product is stretched 1% to 3% in the rolling direction to eliminate residual quenching stresses.
  • the tensile strength of the alloy of the present invention is relatively insensitive to quench rate. Thus its superior strength levels are maintained in plate products of substantial thickness.
  • This property of the alloy produced according to the present invention results from the use of zirconium instead of chromium as the grain refining element. Chromium is used for most other 7000 series alloys and results in substantial decreases in strength for section thicknesses over about 7.6 cm (3 inches) whereas the alloy produced according to the present invention decrease only moderately in strength even when produced in section thicknesses well over 7.6 cm.
  • the presently preferred method to artificially age the product produced from the alloy produced according to the present invention is to use a two step aging procedure.
  • the alloy is preferably first aged at an intermediate temperature on the order of 121 °C. for a period of from about 4 to 48 hours. It should be noted that the first aging step can be modified or even possibly eliminated. For example, data accumulated to date indicates that the alloy can be aged during the first stage at temperatures ranging from 107°C. to 135°C.
  • the second stage aging treatment is conducted at a temperature that is above the aging temperature employed during the first stage.
  • the second stage aging is preferably conducted in the range of from 154°C. to 163°C. until the alloy reaches peak strength.
  • peak strength it is meant a strength at or near the maximum strength of the alloy.
  • the aging time will range from about 3 to about 5 hours. If the second stage aging is conducted at 154°C., the aging time will range from about 6 to about 12 hours.
  • the second stage aging can also be conducted at temperatures in an expanded range of from 149°C. to 171°C. until peak strength is achieved.
  • the aging time must be adjusted upwardly and for temperatures toward the upper end of the foregoing range, the aging time must be adjusted downwardly.
  • the formula below may be used to determine the preferred second stage aging time (t T ) foraging temperatures other than 163°C. This formula will provide an aging time for a given temperature within the range of 149°C. to 171°C. that is equivalent to the second stage aging time for the aging temperature of 163°C. as set forth in the preceding paragraph.
  • the formula is: wherein t T is the time for which the product of the present invention is aged during the second stage aging at a temperate T other than 163°C. to achieve peak strength,
  • the factor Y is derived from the graph of Figure 1 which is a loglinear graph of the Y factor versus aging temperature. For example, if it were desired to conduct the second stage aging at a temperature of 156°C., the factor Y would be about 0.5; and if it were desired to age at a temperature of 170°C., the factor Y would be about 2. It should also be realized that the aging time (t T ) calculated from the above formula can be varied up to about 3 hours and still achieve the peak strength properties in accordance with the present invention. For example, for example stage aging temperatures near the upper limit of the expanded range, the variation from t T is preferably no more than about ⁇ 1/2 hour; however, at the lower end of the expanded range, t T can be varied up to about ⁇ 3 hours.
  • More than fifty ingots of the alloy produced according to the present invention were formulated in accordance with conventional procedures. These ingots had a nominal composition of 6.4% zinc, 2.35% magnesium, 2.2% copper, 0.11% zirconium, 0.07% iron, 0.05% silicon, ⁇ 0.01 % manganese, 0.01 % chromium, 0.02% titanium, and a total of ⁇ 0.03% of other trace elements, the balance of the alloy being aluminum.
  • the ingots were rectangular in shape and had thicknesses between 41 and 61 centimeters.
  • the ingots were scalped, homogenized at about 471 °C., and hot rolled to plate thicknesses varying from .953 cm. to about 3.8 cm.
  • the 7075 alloy had a nominal composition of 5.6% zinc, 2.5% magnesium, 1.6% copper, 0.2% iron and 0.15% silicon, the balance of the alloy being aluminum and small amounts of other extraneous elements.
  • the 7178 alloy had a nominal composition of 6.8% zinc, 2.7% magnesium, 2.0% copper, 0.2% chromium, 0.05% manganese, 0.2% iron and 0.15% silicon, the balance of the alloy being aluminum and small amounts of other extraneous elements.
  • the 7050 alloy had a nominal composition of 6.2% zinc, 2.25% magnesium, 2.3% copper, 0.12% zirconium, 0.09% iron, 0.07% silicon, 0.01% chromium, 0.02% titanium, the balance of the alloy being aluminum and small amounts of other extraneous elements.
  • K a pp The fracture toughness parameter (K a pp) is related to the stress required to fracture a flat panel containing a crack oriented normal to the stressing direction and is determined from the following formula:
  • the data for the fatigue crack growth rate comparisons was taken from data developed from precracked, single edge notched panels. The panels were cyclically stressed in laboratory air in a direction normal to the orientation of the fatigue crack. The minimum to maximum stress ratio (R) for these tests was 0.06.
  • Fatigue crack growth rates (da/dN) were determined as a function of the cyclic stress intensity parameter (AK) applied to the precracked specimens.
  • the parameter AK(MPa4m) is a function of the cyclic fatigue stress ( ⁇ ) applied to the panel, the stress ratio (R), the crack length and the panel dimensions. Fatigue comparisons were made by noting the cyclic stress intensity (AK) required to propagate the fatigue crack at a rate of 0.185 microns/cycle for each of the alloys.
  • the AK level required to provide a crack growth rate of 0.185 microns/cycle for the 7075-T651 alloy was about 11 MPa ⁇ m; for the alloy of the present invention, 12 MPa ⁇ m; for the 7178 alloy 9.0 MPa ⁇ m; and for the 7050 alloy, 12 MPavm.
  • the bar graphs in Figure 2 show that the alloy produced according to the present invention has strength, fracture toughness and fatigue properties that are 10 to 15% better than the 7075-T651 base line alloy.
  • the 7050-T7651 alloy has fracture toughness and fatigue properties similar to that of the invention alloy, however, the compression yield strength of the 7050-T7651 alloy is not only below that of the alloy of the present invention but is also slightly below that of the base line alloy 7075-T651.
  • the fracture toughness and fatigue crack growth rate properties of the invention alloy are substantially improved over those of the 7178-T651 alloy.
  • Example I The procedures of Example I were employed to produce a plate product from typical ingots of the alloy produced according to the present invention. After initially artificially aging the products for about 24 hours at about 121 °C., the products produced from the alloy of the present invention were subjected to a second stage aging step at 163°C. for varying amounts of time ranging from 0 to 24 hours. The alloys had the same nominal composition as the alloys produced according to the present invention shown in Example I. Specimens taken from the products were then tested for longitudinal yield strength using conventional procedures. The resulting typical yield strengths versus aging time are plotted in Figure 3. Graph A indicates the strength values obtained from the plate product of the invention.
  • the invention alloy produced achieves and maintains peak strength after additional aging at 163°C. for about 3 to 5 hours. To the contrary, as the 7075 and 7178 plates are exposed to the 163°C. second stage aging treatment, their strength immediately begins to decrease. It is also observed that when the alloy produced according to the present invention is overaged significantly, on the order of 15 to 25 hours, its strength falls below its peak or maximum strength. At these significantly overaged tempers, however, the alloy produced according to the present invention shows significant improvements in short transverse stress-corrosion resistance and exfoliation resistance.
  • the fracture toughness for the product produced from the alloy produced according to the present invention is shown in graph D of Figure 4, the fracture toughness for the 7075-T651 alloy by graph E, and the fracture toughness for the 7178-T651 alloy by graph F.
  • the alloy of the present invention exhibits betterfracture toughness than alloy 7075-T651 and much improved toughness compared to alloy 7178-T651.
  • an alloy having the composition of the alloy produced according to the present invention was formed into plate products of varying thickness in accordance with the procedure set forth in Example I, with the exception that the hot rolling temperatures were not sufficiently high to prevent excessive recrystallization in the plate products. It was determined that approximately 75 volume percent of the alloy was recrystallized.
  • the room temperature fracture toughness data for these substantially recrystallized plates of the alloy are plotted versus plate thickness in graph G of Figure 4. As will be observed, the fracture toughness properties of the invention alloy, when substantially recrystallized, fall to approximately the levels of the 7178-T651 alloy. As a consequence, it is important that the alloy produced according to the present invention be hot rolled in a manner that will prevent substantial recrystallization.
  • the volume percent recrystallized was determined for this Example by the point count method on photomicrographs (100x magnification) of a full thickness sample.
  • the alloy of the present invention for which fracture toughness data is presented in graph E of Figure 4 was only about 17% recrystallized, while the alloy for which fracture toughness data is presented in graph H was about 75% recrystallized. From this, it is apparent that an alloy produced according to the present invention must be substantially unrecrystallized in order to provide fracture toughness properties that are better than the prior art alloys.
  • the fatigue crack growth rate (da/dN) properties of the alloy produced according to the present invention are improved over other commercial alloys having similar strength characteristics, namely the 7075-T651 and 7178-T651 alloys.
  • Four production lots of plate material of the alloy produced according to the present invention were prepared in accordance with the general procedure set forth in Example I.
  • nine production lots of 7075-T651 alloy plate and two production lots of 7178-T651 alloy plate were procured.
  • fatigue crack growth rate tests were conducted on precracked single edge notched panels produced from the production lots of each of the alloys.
  • FIG. 5 is a plot of the mean values of the crack growth rates (da/dN) in microns per cycle versus the cyclic stress intensity parameter (AK) for each of the alloys.
  • Curve H represents the crack growth rates for 7178-T651 alloy, curve I for 7075-T651 alloy, and curve J for the alloy produced according to the present invention.
  • the alloy of the present invention has superior fatigue crack growth rate properties at each stress intensity level examined when compared with the 7178-T651 and 7075-T651 alloys.
  • the alloy produced according to the present invention has a superior combination of strength, fracture toughness and fatigue resistance when compared to the prior art alloys typified by 7075-T651, 7178-T651 and 7050-T7651.
  • Other tests conducted on the alloy produced according to the present invention and comparable 7075-T651 and 7178-T651 alloys also indicate that the stress corrosion resistance and exfoliation corrosion resistance of the alloy of the present invention are approximately equivalent to the corrosion resistance properties of alloy 7075-T651, and thus can be employed for the same applications, such as wing panels and the like.

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Claims (1)

1. Verfahren zur Herstellung eines plattenförmigen Aluminium-Produkts für eine obere Tragflächenhaut eines Flugzeugs, gekennzeichnet durch die Stufen:
(1) Bereitstellung eines Legierungskörpers bestehend aus einer Aluminium-Legierung der 7000-Reihe vom Aluminium-Zink-Magnesium-Kupfer-Zircon-Typ mit einer Zusammensetzung von (Gew.-% der Gesamtlegierung):
Figure imgb0005
(2) Heißbearbeiten des Legierungskörpers durch Heißwalzen zur Herstellung des plattenförmigen Legierungsprodukts, wobei das Heißwalzen positiv kontrolliertwird, indem absichtlich die Temperatur des Legierungskörpers auf einem genügend hohen Niveau gehalten wird, wodurch die Mikrostruktur des plattenförmigen Legierungsprodukts weniger als 50 Vol.-% rekristallisiert ist;
(3) Unterwerfung des plattenförmigen Legierungsprodukts einer Lösungsglühbehandlung und einem Abschrecken;
(4) Strecken des plattenförmigen Legierungsprodukts um 1,5 bis 3% in Walzrichtung; und
(5) Unterwerfung des plattenförmigen Legierungsprodukts einer künstlichen Alterungsbehandlung bei erhöhter Temperatur bis die Spitzenfestigkeit erreicht wird,
wobei das genannte Verfahren in der Weise durchgeführt wird, daß das hergestellte plattenförmige Legierungsprodukt mit einer Dicke von 0,953 cm (3/8 in) bis 3,8 cm (1,5 in) und einer Mikrostruktur von weniger als 50 Vol.-% rekristallisiert, bestimmt unter Verwendung eines Probekörpers des plattenförmigen Legierungsprodukts mit voller Dicke, die folgende Kombination von Kompressionsfließfestigkeit, Bruchzähigkeit und Ermüdungseigenschaften besitzt:
(1) eine Festigkeit, gemessen als minimale Kompressionsfließfestigkeit, Fa, von 524 MPa;
(2) eine Bruchzähigkeit, gemessen durch die mittlere Bruchzähigkeit, Kopp für eine Probekörperdicke von 1 cm, von 77 MPa√m; und
(3) ein Ermüdungsverhalten, gemessen durch die mittlere cyclische Spannungsintensität, AK, bei einem Spannungsverhältnis (R) von 0,06 und in Laboratoriumsluft, von 12 MPa√m, das erforderlich ist, um eine Rißwachstumsgeschwindigkeit (da/dN) von 0,185 µm/Zyklus zu erzeugen.
EP79901364A 1978-09-29 1979-09-24 Verfahren zur herstellung von aluminium-legierungen Expired - Lifetime EP0020505B2 (de)

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
US947089 1978-09-29
US05/947,089 US4305763A (en) 1978-09-29 1978-09-29 Method of producing an aluminum alloy product

Publications (4)

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EP0020505A1 EP0020505A1 (de) 1981-01-07
EP0020505A4 EP0020505A4 (de) 1981-02-04
EP0020505B1 EP0020505B1 (de) 1984-05-30
EP0020505B2 true EP0020505B2 (de) 1993-07-14

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US (2) US4305763A (de)
EP (1) EP0020505B2 (de)
JP (1) JPS6317901B2 (de)
DE (1) DE2953182C3 (de)
GB (1) GB2052558B (de)
SE (1) SE447128B (de)
WO (1) WO1980000711A1 (de)

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USRE34008E (en) 1992-07-28
EP0020505A4 (de) 1981-02-04
DE2953182A1 (en) 1980-12-04
GB2052558A (en) 1981-01-28
JPS55500767A (de) 1980-10-09
DE2953182C2 (en) 1994-09-29
EP0020505A1 (de) 1981-01-07
SE447128B (sv) 1986-10-27
JPS6317901B2 (de) 1988-04-15
GB2052558B (en) 1982-12-08
US4305763A (en) 1981-12-15
SE8003997L (sv) 1980-05-29
DE2953182C3 (de) 1994-09-29
EP0020505B1 (de) 1984-05-30

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